Heat resistant alloys

ABSTRACT

The development of a new family of alloys, having exceptional stability and high strength at temperatures beyond three-fourths of their absolute melting points, has been made possible by the discovery of an entirely new mechanism wherein supersaturation of the matrix occurs during heating as opposed to supersaturation during cooling, such as with conventional age hardening type alloys.

United States Patent Ghosh 1 Dec. 9, 1975 [5 HEAT RESISTANT ALLOYS3,174,851 3/1965 Buel'llfil' 75/170 [75] Inventor: Subrata Ghosh,Southfield, Mich. OTHER PUBLICATIONS [73 Assignee; Chryskr corporafion,Highland Metals Handbook, 1948 Edition, American Society park, i forMetals, pp. 1158, 1163, 1164, 1206, 1220, 1221, 1243, and 1245. [22]Filed: May 20, I968 [21] Appl. No.: 730,226 Primary Examir1erC. LovellAttorney, Agent, or FirmTalburtt & Baldwin 52 U.S.Cl. 7 7 4 2; 148 32. l1 5/1 0 1 8/3 72/1751 [57] ABSTRACT 51 int. Cl. C22c 19/00 Thedevelopment of a new family of y having 53 Fi f Search 0 75/170, 17],135, 147, ceptional stability and high strength at temperatures 75/122;148/32.5, 32 beyond three-fourths of their absolute melting points, hasbeen made possible by the discovery of an entirely [56] References cunew mechanism wherein supersaturation of the matrix UNITED STATESPATENTS occurs during heating as opposed to supersaturation duringcooling, such as with conventional age harden- 2,570,193 10/1951B|eber..... ing type alloys 2,910,356 10/1959 Grala 3,021,211 2/1962Flinn 75/170 1 Claim, 8 Drawing Figures U.S. Patent Dec. 9, 1975 Sheet 1of 3 K-FL 11/1302 l NIL/7!.

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IT TOW/145') US. Patent Dec.9, 1975 Sheet 2 of3 3,925,071

U.S. Patent Dec. 9, 1975 Sheet 3 of3 3,925,071

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HEAT RESISTANT ALLOYS BACKGROUND OF THE INVENTION The advent of the gasturbines and space exploration vehicles has imposed an ever-increasingdemand for the development of superior high temperature resistant alloysfor prolonged use as structural members. To meet this demand,metallurgists have constantly improved on their earlier achievements bynew discoveries and innovations based on the fundamental principles ofalloy strengthening in view of heat resistance.

The fundamental definition of heat resistance in the solid state is theability to resist plastic deformation and failure under the action ofstresses at high temperatures over a period of time. Plasticdeformation, or dislocation glide, is inhibited by inherent structuralmeans which interfere with the mobility of dislocations. To provide forthese structural aspects, the metallurgists have taken recourse to themethods of: (1) solid solution, wherein the base metal is alloyed withother elements up to its saturation limit; (2) dispersion, wherein asecond insoluble phase is uniformly dispersed within a pure metal or asolid solution matrix; (3) carbide for mation, wherein highly stablecarbide structures are formed by the alloy components and distributedwithin the matrix in critical form and areas; and (4) precipitation,wherein a supersaturated solid solution is made, under controlledconditions, to reject the excess solutes in the form of second phaseprecipitates within the matrix in critical form and areas. The objectiveof providing interference to dislocation mobility may be achieved byeither one or careful combination of the above methods, depending on thespecific requirements.

While these heretofore employed strengthening mechanisms have been usedboth singularly and in combination to provide the numerous hightemperature alloys commercially available today, each is accompanied bylimitations which resulted in the usetemperature of the alloy beingrestricted within onehalf to three-fourths of its absolute meltingtemperature (0.5-0.75 Tm). Thus, solid solution strengthening is limitedto a narrow group of element combinations and their solid solubilityranges with respect to a particular base element. Dispersionstrengthening, on the other hand, is difiicult to achieve due first tothe lack of wide compatibility ranges between metallic matrices andbasically non-metallic dispersion phases and secondly, due to the needfor complicated powder metallurgy or expensive electro-depositionprocesses to introduce the dispersion phase. Carbide strengthening islimited to the increasing degree of brittleness introduced with higherpercentages of carbides, the limited stability of the particularcarbides formed in the alloy, as well as the compatibility ranges of thecarbide structure with the matrix of the alloy.

One of the most important mechanisms employed by the metallurgist isprecipitation hardening. The atoms of a pure, solid, crystalline elementarrange themselves in a specific order in space which give rise to whatis known as their unit cell. This basic cell structure may accept otheratoms either in place of or in-between its own atoms resulting in asolid solution. Whether this acceptance of foreign" atoms takes place atrandom or to a certain limit, depends basically on the types of thehost" and foreign atoms and the temperature. Usually, in combinationsother than a stoichiometric chemical compound, immiscible mixtures or acontinuous solid solution, this limit of acceptance or solid solubilitydecreases as the temperature is lowered below the solidus temperature.Thus, at a reduced temperature the system becomes supersaturated withthe foreign solute atoms and assumes a state of lower free energy byrejecting some of these atoms from its structure and consequentlybecomes more stable. The rejected atoms, in turn, form a structuralgroup of their own, with or without some of the host atoms, andgradually, through a step by step transitional atomic configura tion,appear in the matrix or host structure as second phase precipitates. Theprecipitates act as anchors to the dislocations and limit their mobilityin the structure. However, due to the very fact that the precipitatesare more stable at lower temperature and begin to dis solve in thematrix as their solid solubility increases with temperature, theanchoring of the dislocation gradually diminishes in strength and beginto yield to their advancing glide under stresses. Thus, thestrengthening effects of the precipitates are eventually lost above thetemperature where their diminishing size and number cannot cope with theincreasing number of dis locations at stress leading to a yieldingmatrix. Consequently, all precipitation or age-hardening typesuperalloys are usable only up to a temperature range of 50-75% of theirabsolute melting points (0.5-0.75 Tm).

Heretofore, the approach which has been used to provide high temperaturealloys is the production of alloy structures which are extremely hardand strong at low temperatures so that despite substantial reduction atelevated temperatures, the alloy is expected to retain sufficientstrength for its intended use. Consequently, these alloys do not readilylend themselves to machin ing and other mechanical working at normaltemperatures.

The present invention concentrates upon this basic weakness in theheretofore designed alloys for use at high temperatures and relates toalloys and method for their production which would provide superior highstrength properties at temperatures up to of their absolute meltingpoint (0.90 Tm).

SU MMARY An object of the present invention, therefore, is to provide anew mechanism for developing resistive strength in an alloy as thetemperature of the alloy increases.

A further object of this invention is to provide new alloys havingsuperior strength at temperatures in the range of from about one-half upto about ninety percent of their absolute melting point (0.9 Tm).

An additional object is to provide relatively inexpensive casting typealloys of low density which are remeltable without sacrificing desirableproperties and which are composed of a minimum amount of costly or heavyelements.

Yet another object is to provide new high temperature alloys having astrengthening phase which is effective at temperatures close to thealloy's melting point, but which disappears on cooling so that the alloycan be readily worked at normal temperatures.

Other objects and advantages of the present invention will becomeapparent from a further reading of the description and the appendedclaims.

The above and other objects and advantages of this invention areobtained by my discovery that elements 3 of a complex alloy may bechosen such that, within specified composition ranges, supersaturationof the matrix takes place during heating as opposed to super saturationduring cooling as in the case of conventional precipitation hardening.

DESCRIPTION OF THE DRAWINGS FIG. I, shows a portion of the phase diagramof the binary nickel-aluminum system;

FIG. 2, is an enlarged view of a portion of the nickelaluminum phasediagram, the portion being indicated generally by the area designated asA" in FIG. I;

FIG. 3, illustrates a portion of the phase diagram of the binary copperaluminum system;

FIG. 4, illustrates a portion of the phase diagram of the binaryaluminum-cobalt system;

FIG. 5, illustrates a portion of the phase diagram of the binarycobalt-tungsten system;

FIG. 6, illustrates a portion of the phase diagram of the binarycopper-zinc system;

FIG. 7, illustrates the phase diagram at a temperature of [382 F. of theternary nickel-aluminum-chromium system; and

FIG. 8, illustrates the phase diagram at a temperature of [832 F. of theternary nickel-aluminum-chromium system.

DESCRIPTION OF THE INVENTION As mentioned above, heretofore knownprecipitation or age hardening type superalloys are only recommended foruse at temperatures below about 75 percent of their absolute meltingpoint. The reason for this being that the precipitates are stable at atemperature below the saturation limit, but gradually dissolve away asthe temperature of the alloy increases. Also. the other strengtheningsources begin to diminish gradually with the rise in temperature. Thus,it is seen that the major strength of the alloy at higher temperaturesdepends on the precipitates and those precipitates gradually becomenon-existent when higher temperatures are encountered.

In contrast with the above situation, I have now discovered that it ispossible to develop alloys in which the dislocation anchoringprecipitates will behave in an exactly opposite manner. Thus, at lowertemperature, when there is no need for critical strengthening, theprecipitates dissolve out; however, at higher temperatures. a reactiontakes place within the aggregate of the host and alloying foreign atomsdue to supersaturation of the matrix. This reaction results in theprecipitation of an additional phase. Accordingly, the general softeningeffects of the high degree of heat energy is resisted by the appearanceof precipitates at the increased temperature.

In utilizing the new strengthening mechanism of this invention todevelop superior high temperature alloys having excellent strengthcharacteristics up to about ninety percent of their melting point.resort is made first to the phase equilibrium diagrams from which aninitial approximation of a base alloy composition may be made in whichthe alloy enters. upon being heated. into a solid state zone containingan additional phase. Thus. the method of the invention for producingalloys which exhibit superior strength properties at elevatedtemperatures comprises selecting first a base element and at least onealloying element which are capable of combining to form an alloy thatundergoes a solid state to solid state phase transformation, andcombining them to form an alloy which, upon heating, undergoes thistransformation so as to form at least one additional phase at or belowthe temperature at which it will be used. It should be understood thatthis invention is not merely limited to alloy systems wherein the alloycomposition is a single phase at room temperature and two phases at thedesired use temperatures. It is only necessary in this invention that atleast one additional solid state phase be formed at or near the desireduse temperature. Accordingly, the strengthening mechanism of thisinvention is found, for example, in those alloys comprising two or moresolid state phases at normal temperature and three or more solid statephases at use temperature, or alloys comprised of one phase when cooledand three or more solid state phases when heated.

Examples of binary systems wherein it is possible to combine elementsaccording to the invention into a1- loys which form at least oneadditional phase upon heating include, but not limited to those shown inthe following table. In this table, it will be seen that a binary alloycomposed of about 72 to 73 atomic percent aluminum and 27 to 28 atomicpercent cobalt will. when heated to a temperature in the range of about[832 to 2012 F. enter a phase region where an additional solid statephase will precipitate. Likewise, it is seen in the magnesium-zincbinary system that there are three distinct composition ranges in anyone of which an alloy can be formed in which precipitation will occurupon the alloy being heated to a temperature in the range of about572-662 F. It should be understood that the composition and temperaturevalues shown in the following table are of a representative value andmay de viate somewhat from actual limits.

TABLE A Atomic Binary System Percent Temperature Range Elements of F. of

A B Element B Precipitation Aluminum Cobalt 27-28 1832-2012 AluminumIron 25-26 1832-2120 Aluminum Magnesium 44-45 572-734 Aluminum Titanium50-51 2372-2660 Aluminum Tungsten 22-26 1832-2372 Antimony Tin 59.5-212-464 Bismuth Lead 67-77 86-360 Cadmium Lidtium 16-25 302-626 CadmiumNickel 17-21 752-932 Cerium Thorium 0.1-15 1238-1526 Chromium Tantalum34-36 2912-3542 Cobalt Antimony 64-67 1 1 12-1652 Cobalt Osmium [-35752-2822 Cobalt Rhenium [-25 752-2732 Cobalt Ruthenium [-33 842-2552Cobalt Tungsten 7-13 1292-1922 Cobalt Vanadium 24-31 1922- l 958 CopperAluminum l6-19.6 1040-2066 Copper Antimony 15-20 752-1022 Copper Cadmium41.5-44.5 572-1022 Copper Gallium 16.3-18.6 1 148-1670 Copper Germanium12-13 1 112-1472 Copper Magnesium 32-34 932-1292 Copper Silicon17.5-18.5 752-1292 Copper Zinc 31.9-38.3 842- 1652 Gold Cadmium 27-34932-1148 Gold Indium 15-21 932-1202 Gold Manganese 46-52 1022-1130Indium Lead 27-31 266-338 Indium Thallium 43-58 68-302 Indium Tin 13-27167-248 Iron Columbium 34-36 1832-2372 Iron Nitrogen 5.7-6.1 842-1202Iron Tungsten 62-63 1832-2000 lron Tungsten 69-70 2600-2800 IronVanadium 43-52 2012-2246 Iron Zinc 69-709 1238-1436 Magnesium Cadmium20-35 212-356 Magnesium Cobalt [-15 392-2012 Magnesium Indium 23-25482-644 TABLE A-continued Atomic Binary System Percent Temperature RangeElements of F. of

A B Element B Precipitation Magnesium Lithium 17-175 572-1076 MagnesiumZinc 49.5-50.5 572-662 Magnesium Zinc 595-605 572-662 Magnesium Zinc84-85 572-662 Manganese Nickel 47-55 1238-1400 Mercury Lead 66-69122-302 Nickel Gallium 40-42 1112-1382 Nickel Germanium 23-25 1832-2102Nickel Molybdenum 28-29 1472-1652 Nickel Molybdenum 34-35 1472-1652Nickel Platinum 55-70 752-1 1 l2 Nickel Silicon 22-23 1652-2012 NickelVanadium 22-27 1 1 12-1832 Nickel Zinc 45.5-51.5 1256-1472 PalladiumLead 38-41 1346-1526 Palladium Zinc 32-34 1202-1292 Palladium Zinc 37-561 1 12-2102 Rhodium Tin 38-405 1652-4352 Silver Aluminum 21-26 392-842Silver Aluminum 24.5-32.5 1 130-1328 Silver Cadmium 384-40 1 1 12-1360Silver Magnesium 74-78 392-878 Silver Mercury 44-45 32-572 SilverPlatinum -18 752-1742 Silver Tin 12-13 392-1292 Silver Tin 23.7-25.5572-896 Silver Zinc 32-40 572-1292 Silver Zinc 58.5-61 752-1220 SilverZinc 67-69 932-1 166 Thallium Cerium 55-56 1112-1652 Titanium Tungsten0.2-9.0 1328-1616 Titanium Uranium 59-67 1292-1634 Cerium Yttrium lYttrium 1) Lanthanum Yttrium 1 Magnesium Yttrium l) Thorium Zinc Cerium10-1 3.5 932-1472 Zinc Cobalt 7-8 842-1040 Zinc Lithium 13-23 149-464Zinc Magnesium 13-195 680-752 Zinc Magnesium 23-27 500-610 1 Existenceof Inverse Precipitation Reaction in these Systems is Established, butComposition and Temperature are not well defined.

From the preceding table, it is seen that the mechanism of thisinvention is applicable to relatively high melting alloys such asiron-tungsten (2600-2800F.), medium temperature melting alloys ascopper-zinc (840-l650F.), and low temperature systems such asbismuth-lead (86-360F.). Thus, it will be understood that this inventionis not limited to alloys of high melting temperature, but instead isconcerned with strengthening of the alloy at temperatures close to themelting point of the alloy, whatever that may be. It will likewise beunderstood that alloys based on this invention may very well be a binaryalloy, but more often than not further alloying will be necessary toachieve the desired temperature strength.

From the foregoing, it is seen that this invention is applicable to manyand varied alloy systems and, accordingly, for purposes of clarity inthe ensuing discussion of the invention, reference will be made only tothe nickel-aluminum binary system. P10. 1, illustrates a portion of thenickel-aluminum binary system containing from about 50 to 100 weightpercent nickel, the balance being aluminum. It will be noted that thealloys containing from approximately 83.0 to 85.5 weight percent nickeland 17 to 14.5 weight percent aluminum are composed of a single solidstate phase which is the intermetallic compound Ni Al at roomtemperature. However, when heated, a solid state to solid statetransformation reaction occurs and the alloy moves into a binary phase.This can be seen more clearly with reference to FIG. 2, wherein thesingle solid state phase region of Ni Al is hatched. Line 1 denotes theleft hand boundry of this single phase region and it is seen that if analloy composition such as A is heated in excess of about 2200F. thetransformation line 1 is crossed and the alloy moves in a zone where anadditional solid state phase of NiAl is found. This means that the alloywhich is essentially percent Ni,-,Al at lower temperatures will begin toprecipitate NiAl at a temperature, which will vary with the compositionof the alloy. Likewise, it is seen that alloy compositions such as B andC when heated to approximately 2500F. will cross transformation line 2in which event an additional solid state phase (nickel solid solution, 7will commence to precipitate. Accordingly, it is seen that by carefulselection of elements, it is possible to form an alloy which resistssoftening at elevated temperatures by virtue of bringing into the alloystructure new dislocation anchoring precipitates. Of course, this maynot mean that the alloy becomes stronger than when it is at lowertemperature; however, its heat resistance in comparison with heretoforeknown alloys is greatly increased since, in contrast to conventionalprecipitation hardened alloys wherein the precipitates dissolve onheating, here the precipitates are actually formed on heating.

As mentioned earlier, the Inverse Precipitation" mechanism of thisinvention can be used in many and varied alloy systems, By way ofillustration reference is made to FIG. 3, showing a portion of thecopperaluminum binary system in which it is seen that an alloy having acomposition in the range of about 7.5 weight percent aluminum-92.5weight percent copper (composition C) to about 9.8 weight percentaluminum-90.2 weight percent copper (composition D) will pass throughline 3 upon heating. and undergo a transition from the single solidstate a phase to the binary solid state a+B phase in which 5 will be aprecipitate.

FIG. 4, illustrates a portion of the aluminum-cobalt binary phasesystem. In this system, alloy compositions of from about 46.6 to 48weight percent cobalt-52 to 53.4 weight percent aluminum will, uponbeing heated to a temperature of about 2020F., pass through thetransformation line 4 and change, as seen with respect to composition E,from the single solid 5 phase to the binary 8+ 6 in which 6 will formthe precipitate.

FIG. 5, illustrates a portion of the cobalt-tungsten binary system andit is seen that alloys having a composition in the range of about 44weight percent tungsten- 56 weight percent cobalt (composition F) toabout 48 weight percent tungsten-52 weight percent cobalt (compositionG) will, when heated to a temperature higher than the transformationlines 5 or 5' undergo a transition from the single solid 1 phase to abinary phase of 7 plus B or y plus 8 in which B or 8 respectively, willform the precipitate. FIG. 6 illustrates a portion of the copper-zincphase system. Here, again alloys having compositions in the range ofabout 32.5 weight percent zinc-67.5 weight percent copper (compositionH) to about 38 weight percent zinc-62 weight percent copper (compositionJ) will cross the transformation line 6 and enter the solid phase B Bregion wherein B will form the precipitate. These illustrations indicatethe feasibility range of the lnverse Precipitation mechanism in theparticular alloy system and development of a particular alloy can beachieved with routine experimentation.

The temperature at which the additional solid phase formation beginswill, as seen from the drawings, depend on the composition of the alloy.For most applications, it is preferable to prepare an alloy compositionin which the additional solid state phase appears when the alloy isheated to a temperature in the range of from about A to of its absolutemelting point (Tm). Absolute melting point being the melting temperatureplus 460 in the Fahrenheit scale (Rankine) or plus 273 in the Centigradescale (Kelvin).

It is apparent from the drawings that many of the alloy compositionranges in which it was possible to obtain high temperature precipitationare rather narrow if only a binary system is considered. Although it istrue that a binary alloy, prepared according to this mechanism may besuitable for certain high temperature application, the critical effectsof adding further alloying elements, particularly for expansion of thehigh temperature precipitation zone as well as additional strengtheningis a logical step to consider. Accordingly, as in conventional alloys,the use of additional elements for alloying is employed. It was foundthat through proper selection of the alloying elements it is possible tocontrol the phase regions so as to increase the range of alloycompositions in which the high temperature precipitation mechanism ofthis invention can be em ployed, as well as, alter the transformationtemperatures for a particular base composition. For example, FIG. 7illustrates the ternary nickel-aluminumchromium phase diagram at atemperature of approximately l382 F. it will be noted that the singlephase region of Ni Al is considerably larger than in the binary systemof FIGS. 1 and 2 up to a critical amount of chromium addition.Accordingly, this means that the range of alloy compositions which willundergo transition, upon heating, to form an additional solid statephase is considerably increased. Thus, an alloy of composition X in FIG.7 is in the lOO percent Ni Al phase region at l382 F. However, when thetemperature of the ternary alloy X of FIG. 7 is raised to about 1832 F.the phase regions shift and the composition X is now in the binary phaseregion of Ni Al and NiAl as seen in FIG. 8. This means that at thistemperature NiAl precipitates will appear in the alloy structure.Similarly, alloys of Y" and Z compositions will precipitate y and a Brespectively as the higher temperature is reached.

Based on the concept as outlined above, a series of high temperaturesuperalloys containing up to 20 elements and comprising about 7 to 17weight percent aluminum and about 65 to 85.5 weight percent nickel havebeen developed starting with a binary nickelaluminum base compositionwhich is then alloyed with other suitable elements. The preferredelements for use in alloying quantities with a base nickel-aluminumsystem are chromium, niobium, carbon, titanium, cobalt, molybdenum,tungsten, tantalum, boron, silicon, vanadium, beryllum, nitorgen, rareearths, yttrium, zirconium, copper, hafnium, rhenium, oxygen, manganeseand iron. These alloying elements in addition to widening thecomposition zone where the high temperature precipitation mechanism ofthis invention is attainable, strengthen the alloy in conventionalfashion such as by solid solution, conventional precipitation, carbidestrengthening, etc. Accordingly, the alloys provided by this inventionderive their high strength from all the conventional strengtheningmechanisms and, in addition, offer exceptional strength at temperaturesclose to their melting points by means of high temperatureprecipitation. In other words, the strengthening at lower temperaturesis derived from the conventional sources of alloy strengthening, and attemperatures when such conventional sources tend to become exhausted,the high temperature precipitation takes place to maintain theirstrength and further their use temperature beyond three-fourth (0.75 Tm)of their melting point. Thus, the new strengthening mechanism of thisinvention supplements and does not merely replace conventionalstrengthening mechanisms.

It should also be recognized that the high temperature precipitationmechanism of this invention makes it feasible to obtain a stabledispersion phase in a casting type alloy. For example, in thenickel-aluminum base alloys, oxygen can be reacted with the NiAlprecipitate to form a stable dispersion of aluminum oxide according tothe equation: 6NiAl (ppt) 30 2Ni Al 2Al 0 The oxygen can be suppliedfrom dissolved oxygen in the alloy or by diffusion from the outside.

The following examples illustrate alloys prepared in accordance withthis invention:

EXAMPLE 1 An alloy according to the present invention was preparedhaving the following nominal analysis:

Nickel 79.4 weight percent Aluminum 134 weight percent Chromium 7.2weight percent This alloy was prepared by simultaneously charging allthe elements into an induction furnace enclosed within a vacuum chamber,and then evacuating the chamber to produce pressure in the furnace ofapproximately 1-10 microns. The furnace was then heated and the chargemelted, forming a molten mass having a temperature between 25002 600 F.The heating was continued until the melt reached a temperature of 3200F. so as to ensure homogenization. After about 5 minutes at thistemperature, the melt was cooled to about 3050 F. and poured into apre-heated mold located inside the vacuum chamber so as to form testbars. The test bar castings were withdrawn from the vacuum chamber abouteight minutes after casting and allowed to cool in air. The test barswere approximately 4 inches long and /2 inch in diameter at the endportions with a reduced A inch diameter center neck portion of l-V2inches.

It was determined that from room temperature to approximately 1560 F.the alloy was the single Ni Al phase and that above 1560 F. NiAl beganto precipitate. At about 2l00 F. the amount of NiAl reached about 15 to20 percent of the matrix. The alloy had a theoretical density of 0.24pounds per cubic inch and a Rockwell C (R hardness of 32-35.

The following data were obtained through tests conducted at roomtemperatures and elevated temperatures on the above prepared alloy.

Stress Rupture Life At: l500F/30,000 psi l 1 Hours 3.0% ElongationEXAMPLE 2 An alloy of the following nominal analysis was prepared as setforth in Example I,

Nickel Aluminum Chromium Titanium 74.8 weight percent 1 L weight percenti2.5 weight percent 1.7 weight percent This alloy had a theoreticaldensity of 0.260 pounds per This alloy had a theoretical density of0.278 pounds per cubic inch and a Rockwell hardness R of 30-35.

The following data were obtained through tests conducted on the aboveprepared alloy.

cubic inch and a Rockwell hardness R of 32-35.

The following data were obtained through tests conducted on the aboveprepared alloy.

Stress Rupture Life At'.

l500F/30.000 psi 59 Hours 6% Elongation Stress Rupture Life At:

i100F/20.00o si meow/20.000 psi 1800F/l5,000 psi wow/10.000 psi TABLE 4Temp. Ult. 02% 7( "F Tensile Stress Yield Elongation 70 l 18 I02 7 i80067 64 3 470 Hours-7 0% Elongation 62 Hours-2.0% Elongation 2i]Hours-5.0% Elongation l36 Hours-2.0% Elongation EXAMPLE 5 An alloy ofthe following nominal analysis was pre 0 pared as set forth in Example1:

Nickel Aluminum Chromium Titanium Cobalt Molybdenum Tantalum NiobiumTungsten Carbon Boron Rare Earths Zirconium 70.42 weight percent 7.82weight percent 9.27 weight percent L82 weight percent 3.47 weightpercent 2.93 weight percent 1.92 weight percent 0.25 weight percent 0.50weight percent 0.50 weight percent 005 weight percent 0.05 weightpercent L00 weight percent This alloy had a theoretical density of 0.275pounds per cubic inch and a Rockwell hardness R of 28-36. Melting ofthis alloy began at 2300F. (2760 in Absolute Rankine Scale). Nodeterioration in properties was noted in this alloy on remelted andre-cast condition. The following data were obtained through testsconducted on the above prepared alloy.

EXAMPLE 3 An alloy of the following nominal analysis was prepared as setforth in EXAMPLE 1:

Nickel 7 l.l5 weight percent Aluminum 9.50 weight percent Chromium 9.65weight percent Titanium 2.54 weight percent Cobalt 3.13 weight percentMolybdenum 3.05 weight percent Tantalum 0.95 weight percent Boron .03weight percent This alloy had a theoretical density of 0.269 pounds percubic inch and a Rockwell hardness R of 30-35.

The following data were obtained through tests conducted on the aboveprepared alloy.

TABLE 3 Temp. Ult. 0.2%

"F Tensile Stress Yield Elongation 70 I23 I l 1 3 I800 65 63 2 StressRupture Life At:

l500F/25,000 psi 1837 Hours 2% Elongation l500F/30,000 psi 1000 Hours 2%Elongation l500F/40,000 psi 410 Hours 1% Elongation l700F/20,000 psi 109Hours 1% Elongation EXAMPLE 4 An alloy of the following nominal analysiswas prepared as set forth in Example 1:

Nickel 71.30 weight percent Aluminum 7.93 weight percent Chromium 9.38weight percent Titanium 1.85 weight percent Cobalt 3.5l weight percentMolybdenum 2.97 weight percent Tantalum 1.94 weight percent Niobium 0.25weight percent Tungsten 0.50 weight percent Carbon 0.30 weight percentBoron 0.05 weight percent Rare Earths 0.05 weight percent Stress RuptureLife At:

1800F/20,000 si l800F/l5.000 psi 2000Fl 7,500 psi TABLE 5 Temp. Ult.0.2%

F Tensile Stress Yield Elongation I37 Hours-7.0% Elongation 548Hours-7.0% Elongation 326 Hours-4.0% Elongation The development of thealloys in the foregoing exam- The validity of the concept that anequilibrium phase transformation in the solid state at high temperaturesmay be utilized to strengthen an alloy for high temperature applicationsis, therefore, proven through these examples. As noted earlier, theapplications of this concept may be wide and varied, depending on therequirement for a particular use, for alloys of different base elements.in any event, if the alloying elements are 1 l chosen as outlined in theforegoing. so that high temperature precipitation is involved, the alloywill outperform other high temperature alloys of similar compositionsbut based only on the conventional strengthening mechanisms.

According to the foregoing it will be apparent that the objectives ofthis invention, namely to develop a superior, remeltable, inexpensivecasting type high temperature resistant alloy with temperaturecapability up to about 0.90 Tm, but containing only a small quantity ofexpensive heavy elements, have been achieved. in a direct comparison ofproperties and attributes of the alloy of Example 5, it was found thatthe cost of raw materials and manufacture of one pound of this alloywill be the lowest of all heretofore known high temperature superalloys;its density will be the lowest of all heretofore known high temperaturesuperalloys, and pound per pound, its strength will be superior to thebest of the heretofore known commercial superalloys. Its temperaturecapability will be about 0.90 Tm as against about 0.75 Tm for commercialalloys.

I Claim:

1. A temperature resistant metallic alloy consisting essentially ofabout 65.0 to 85.5 weight per cent nickel and about 7 to 17 weight percent aluminum which combines to form an alloy composition, said alloycomposition having a matrix which is a stable solid solution when at atemperature less than about half of the absolute melting point of saidalloy, said matrix when at a temperature above about half and withintwo-thirds of the absolute melting point of said alloy changing fromsaid solid solution to a supersaturated solid solution such that saidmatrix due to said supersaturation above said temperature of about halfof said absolute melting point forms at least one additional solid statephase in precipitate form to provide strengthening of the alloy matrixabove said temperature of about half of said absolute melting point,said alloy including a stable aluminum oxide dispersion phase which willremain upon cooling, said dispersion phase being formed by reactinggaseous oxygen with said additional solid state phase,

1. A TEMPERATURE RESISTANT METALLIC ALLOY CONSISTING ESSENTIALLY OFABOUT 65.0 TO 85.5 WEIGHT PER CENT NICKEL AND ABOUT 7 TO 17 WEIGHT PERCENT ALUMINUM WHICH COMBINES TO FORM AN ALLOY COMPOSITION, SAID ALLOYCOMPOSITION HAVING A MATRIX WHICH IS A STABLE SOLID SOLUTION WHEN AT ATEMPERATURE LESS THAN ABOUT HALF OF THE ABSOLUTE MELTING POINT OF SAIDALLOY, SAID MATRIX WHEN AT A TEMPERATURE ABOVE ABOUT HALF AND WITHINTWO-THIRDS OF THE ABSOLUTE MELTING POINT OF SAID ALLOY CHANGING FROMSAID SOLID SOLUTION TO A SUPERSATURATED SOLID SOLUTION SUCH THAT SAIDMATRIX DUE TO SAID SUPERSATURATUON ABOVE SAID TEMPERATURE OF ABOUT HALFOF SAID ABSOLUTE MELTING POINT FORMS AT LEAST ONE ADDITIONAL SOLID STATEPHASE IN PRECIPITATE FORM TO PROVIDE STRENGTHENING OF THE ALLOY MATRIXABOVE SAID TEMPERATURE OF ABOUT HALF OF SAID ABSOLUTEE MELTING POINT,SAID ALLOY INCLUDING A STABLE ALUMINUM OXIDE DISPERSION PHASE WHICH WILLREMAIN UPON COOLING, SAID DISPERSION PHASE BEING FORMED BY REACTINGGASEOUS OXYGEN WITH SAID ADDITIONAL SOLID STAGE PHASE.